Enhancing Surface Properties of Cu-Fe-Cr Alloys through Laser Cladding: The Role of Mo and B₄C Additives

Abstract

Laser cladding is a powerful surface treatment technique that can significantly enhance the properties of metal alloys. This study delves into the liquid phase separation behavior of Cu-Fe-Cr alloys under the rapid solidification conditions inherent in laser cladding and evaluates the influence of 4% Mo and 2% B₄C additions on the resulting alloy characteristics. The intensive undercooling characteristic of the laser cladding process facilitates the alloy’s entry into the liquid-phase immiscibility gap, prompting pronounced phase separation. Our investigation reveals the emergence of Fe-rich regions, exhibiting a variety of shapes, set against a continuous Cu-rich matrix. The incorporation of Mo and B₄C was found to modulate the mixing enthalpy and entropy, thereby refining the phase distribution: Mo was observed to prevent the agglomeration of Fe cores, resulting in a dispersion of isolated Fe cores throughout the Cu-rich matrix, while B₄C promoted a more uniform compositional distribution. This study further enumerates the enhancements in microhardness, wear resistance, and magnetic properties of the alloys. Notably, the Cu-Fe-Cr-Mo-B₄C alloy demonstrated a microhardness exceeding 600 HV, a low coefficient of friction around 0.15, high saturation magnetization, and reduced coercivity. These results underscore the efficacy of laser cladding in tailoring the microstructure and properties of Cu-Fe alloys, providing insights for the controlled manipulation of phase separation to optimize surface characteristics for engineering applications.

1. Introduction

When the undercooling is greater than the liquid-phase immiscible gap, the liquid Cu-Fe alloy is separated into Fe-rich and Cu-rich regions. This distinctive characteristic results in Cu-Fe alloys displaying a range of unique properties, and they have significant applications in multiple domains, including magnetoresistive materials, advanced bearings, power transmission engineering, packaging, and others.

However, conventional methods such as casting can lead to macro segregation in Cu-Fe alloys due to the density difference between the two. To mitigate the adverse effects of macroscopic segregation on performance, various approaches have been suggested to achieve immiscible alloys with a uniform microstructure. Wu analyzed the spherical macroscopic segregation patterns of various core–shell structures that emerged during the solidification of Cu-Fe alloys within a microgravity environment [1]. It was observed that as the cooling rate lowered, the duration of liquid phase separation was prolonged, resulting in a greater susceptibility to form into core–shell structures that were influenced by Marangoni convection and surface segregation. Luo examined the rapid solidification of the Fe₄₈Cu₄₈Si₄ immiscible alloy under microgravity conditions. He noted that the two layers of nucleoshells represent the ultimate configuration for the liquid phase separation [2]. Numerous researchers have employed the microgravity approach to create liquid-phase immiscible alloys, including Wang [3,4], Xia [5], and Zhao [6]. Electromagnetic levitation, in addition to the microgravity method, is also utilized to prepare immiscible alloys with a homogeneous structure, like Kobayashi [7], Watanabe [8], Cao [9], and Lin [10]. Xia asserts that greater undercooling significantly enhances the microstructure of the perovskite Fe-Cu alloy and highlights the possibility of achieving almost segregation-free solidification with sufficiently large undercooling [11]. Bai investigated the sub-steady liquid phase separation and rapid solidification behavior of Co₄₀Fe₄₀Cu₂₀ alloy by thermal differential analysis, revealing that the undercooling, cooling rate, convection, and surface tension difference between the two separated phases have an important influence on the coalescence and segregation of the separated phases [12].

Laser cladding (LC) has undergone rapid development as an additive manufacturing process in recent years. It is characterized by high undercooling and can therefore serve as one of the potential methods for the production of miscible Cu-Fe alloys. A variety of immiscible alloys have been prepared using the LC method, including Cu₉₅Fe₅ alloy [13], Cu-Ti-Ni alloy [14], and [Cu_0.6(FeCrC)_0.4]_100−xSi_x alloy [15]. Additionally, Zhou achieved Fep/Cu-Cup/Fe duplex alloys [16] and Cu-Fe-Cr-Si-C immiscible alloys [17] and found that Si can enhance the wear resistance of Cu-Fe alloys. Dai examined how scanning speed during laser cladding affects the arrangement of Cu-Fe composite coatings. He found that higher laser scanning speeds cause dynamic undercooling and increased solute trapping of Cu during Fe-rich dendritic refinement [18]. Zhou examined the impact of Fe content on the properties of Cu-Fe alloy and pointed out that the average microhardness of the composite coating gradually increased with the increasing Fe content [19]. Dai investigated the microscopic morphology of the Cu-Fe composite coating using Fe and Cu as substrates, respectively. The Fe-based composite coating consists of an Fe-rich layer with greater microhardness and a Cu-rich layer with lower microhardness. Meanwhile, the Cu-based composite coating is composed of an Fe-rich intermediate layer with higher microhardness and an upper and lower Cu-rich layer with lower microhardness [20].

The process parameters and chemical composition greatly affect the microstructure and properties of Cu-Fe composite coatings. Numerous studies have focused on elucidating the principles of liquid phase separation, the effects of various solidification conditions on liquid phase separation, and methods for reducing liquid phase immiscible alloy segregation. A satisfactory microstructure and properties can be achieved by controlling the solidification of the alloy. Despite obtaining Cu-Fe alloys with a good microstructure, meeting performance requirements beyond the composition limitations is challenging. To achieve this, it is necessary to enrich the number of components of the Cu-Fe alloys, as demonstrated by Wu or Zhou et al. [6,14,17]. However, there are not enough similar studies available. Molybdenum (Mo), one of the important elements in the steel industry, improves the strength [21], wear resistance [22], and corrosion resistance [23,24] and is widely used in electronic devices such as electronic tubes, transistors, and rectifiers. Taking inspiration from the preceding research, we added a small quantity of Mo to a Cu-Fe-Cr alloy using a high degree of LC undercooling, and we successfully prepared the Cu-Fe-Cr-Mo alloy via LC to obtain a magnetically immiscible alloy with higher hardness and better wear resistance. We also optimized the liquid phase separation of the Cu-Fe-Cr-Mo alloy by adding B₄C, taking into account the impact of B and C on the liquid phase separation.

2. Experimental Design

The matrix comprised AISI 1045 steel, while the LC material comprised a mixture of RCF103 commercial Fe-based alloy powder, pure Cu powder (99.9%), and pure Mo powder (99.9%). The chemical composition of RCF103 is listed in

Table 1. Chemical composition of RCF103 (wt.%).

Components	C	Si	Mn	B	Cr	Ni	Mo	Nb	Fe
Powder	0.07	1.1	0.4	0.23	15.2	5.1	1.0	0.31	Other

, while the powder formulations for various experiments are tabulated in

Table 2. Proportions of experimental powders (wt.%).

Powder No.	RCF103	Cu	Mo	B4C
Powder1	43	57	0	0
Powder2	41	55	4	0
Powder3	40	54	4	2

. RCF103, Cu, Mo, and B₄C were ball-milled using corundum balls (with a particle size of 3–5 mm) as a mixing medium. The mixing process lasted for 500 g/3 h.

The distribution of particle sizes in RCF103 was between 50 and 100 μm, while both Cu and Mo powders had an average size of 100 μm. Based on earlier research by the authors [25,26], we determined the following LC parameters for RCF103: laser power—750 W; laser scanning speed—8 mm/s; diameter of the laser spot—2 mm; overlap—50%; powder feed—11.2 g/min; and the powder nozzle was set perpendicular to the substrate surface. The composite powder is injected into the melt pool using argon gas, resulting in the formation of cladding. The process of laser cladding (LC) was executed on a metal deposition machine (Genertec Shenyang Machine Tool Co. Ltd, Shenyang, China) that was converted from a 3-axis Computer Numerical Control (CNC) machining center. The laser emitter (Raycham, Nanjing, China), equipped with a corresponding powder delivery system, has replaced the Z-axis of the machine, as depicted in Figure 1a. The machine is equipped with a laser type IPG1000, capable of emitting continuous laser light with maximum power of 1000 W and a wavelength of 1070 nm. Argon is used as the powder feeding gas and protective gas to prevent oxidation of the cladding layer.

The laser scanning path and cladding formation process are illustrated in Figure 1b. The chosen process parameters enable the production of deposit blocks with satisfactory internal solidification and a notable absence of defects (namely, porosity, and cracks).

The microscopic structure of the coating was visualized using an ULTRA PLUS field emission scanning electron microscope (Zeiss, Heidenheim, German), and elemental identification of the feature location was achieved using energy dispersive spectroscopy (EDS, Zeiss, Heidenheim, German). The phases of different alloy compositions were examined using an X-ray diffractometer (XRD, Aolong, Dandong, China) of Along brand model AL-Y3000. The 2θ angle range of 20°–100° with a step of 0.09°/s was selected for the detection. The physical topography of the cladding surface was measured using the OLS4100 confocal laser scanning microscope (CLSM, Tokyo, Japan) manufactured by Olympus Corporation. During the measurement, a 405 nm semiconductor laser zoom light source and a white LED zoom light source were utilized. Microhardness is assessed through a Vickers hardness tester (Hengyi Precision Instrument Co. Ltd., Shanghai, China), utilizing a 10N load and a 15s holding time. The test is conducted randomly across the surface and repeated 10 times, with the average value of the 10 tests taken as the quantitative measure of microhardness for each respective sample. Wear resistance of cladding surfaces is evaluated using an MFT4000 surface performance tester (Huahui Instrument Technology Co. Ltd., Lanzhou, China). The grinding ball material utilized is SiN. The applied load force for the wear test amounts to 100 N, while the test duration is set to 30 min. Moreover, the travel distance is set to 20mm with a corresponding speed of 240 mm/s.

The present study also calculated the entropy of mixing and enthalpy of mixing of the designed alloy system during solidification using JMatPro v7.0. The calculations were carried out within the General Steel module, setting the initial and ending temperatures in the thermodynamics to be 293.15 K and 2000 K, respectively, in steps of 10 K.

3. Experimental Results

3.1. Macro Phase Separation

Figure 2 illustrates the surface structure of the composite alloys with various compositions, suggesting that the Cu-Fe alloys share comparable traits of liquid phase separation with previous research [27,28]. In Figure 2a, the Fe-rich area appears bar-shaped in the direction of laser scanning. Additionally, there are free-state Fe cores located around the Fe-rich region. These cores originate from Fe-rich droplets that precipitate during liquid separation and do not amalgamate with larger Fe-rich droplets prior to solidification. In Figure 2b, the introduction of Mo disturbs the distribution of Fe-rich areas, causing a significant number of Fe cores to be dispersed throughout the Cu-rich matrix. Moreover, the surface of the Fe-rich region is much flatter compared to that of Figure 2a. In Figure 2c, the inclusion of B₄C reorganizes the interrupted Fe-rich area, while the independent Fe cores in the Cu-rich matrix are also combined into more sizeable Fe cores. The inclusion of B₄C also brings about a more uniform distribution of Fe-rich areas in the Cu-rich matrix.

Figure 3 illustrates the distribution of Fe cores in Cu-rich substrates of Cu-Fe-Cr-Mo and Cu-Fe-Cr-Mo-B₄C. The Fe core size in the Cu-Fe-Cr-Mo-B₄C alloy in Figure 3b appears significantly smaller compared to the Cu-Fe-Cr-Mo alloy in Figure 3c. However, the density of Fe cores shown in Figure 3b is higher than that in Figure 3a, suggesting that the addition of Mo may favor the formation of Fe cores from the melt. The sizes of 300 randomly selected Fe cores in Figure 3a–c were counted, and the resultant normal distribution of Fe core sizes is presented in Figure 3d–f.

The data presented in Figure 3d,e illustrate that the Fe core sizes in the Cu-rich matrix in both the Cu-Fe-Cr alloy and Cu-Fe-Cr-Mo alloy are concentrated in the range of 2–6 μm. Conversely, the result in Figure 3f shows that the Fe core sizes of the Cu-Fe-Cr-Mo-B₄C alloy are predominantly distributed in the range of 0–4 μm, indicating that B₄C effectively reduces the Fe core sizes. Although there are slight variations in the size distribution of Fe cores between Cu-Fe-Cr and Cu-Fe-Cr-Mo alloys, these can be observed in Figure 3d,e. Compared to the Cu-Fe-Cr-Mo alloy, the Cu-Fe-Cr alloy exhibits a significantly greater number of Fe cores with sizes below 2 μm and within the range of 6 to 8 μm on its Cu-rich substrate. Additionally, the Cu-rich matrix of the Cu-Fe-Cr alloy does not contain Fe cores exceeding 12 μm in dimension. The obtained results indicate that while Mo does not lead to remarkable alternations in the size distribution of Fe cores, it does diminish the amount of Fe cores that are initially either too diminutive or too huge in the Cu-Fe-Cr alloy. This outcome contributes to mitigating the macroscopic segregation in Fe-rich areas from the Cu-rich matrix and fosters the consistent dispersion of Fe cores within the Cu-rich matrix.

3.2. XRD Results of Cu-Fe Alloy with Added Mo and B₄C

Figure 4 shows the XRD test outcomes of Cu-Fe compounds with varying compositions. The outcomes signify that the main diffraction peaks of all three compounds are Cu, Fe, and Cr. For the Cu-Fe-Cr compound, the recognized phases are Cu_0.81Ni_0.19 and Fe_0.64Ni_0.36. The presence of Ni, which comes from RCF103 powder, suggests that there are solid solutions indicating that Ni can be dissolved in both Cu and Fe to some extent. The Cr_0.19Fe_0.7Ni_0.11 phase, showing similarity to the 2θ angle, provides assurance of the existence of a composite solid solution comprising Cr, Fe, and Ni. The Ni content in this phase is lesser in comparison to that found in the Cu and Fe phases, pointing towards an inclination of Ni towards the Cu and Fe phases.

The XRD results of the Cu-Fe-Cr-Mo alloy reveal peaks that correspond to pure copper, indicating that Cu exists as a discrete phase and likely forms a matrix embedded with other elements. In comparison to Mo, Cu exhibits a lower tendency to form compounds with Fe and Cr, which accords with this observation. The identification of a (Fe, Ni) phase suggests that Fe and Ni also form solid solutions. The lack of prominent Mo peaks within the main phase suggests that Mo may be either highly dispersed within Fe-rich regions or its quantity is beneath the threshold of detection for significant peak formation. Nevertheless, the Mo peaks with low intensity that were observed at 2θ angles of 58.608° and 73.682° could indicate the presence of trace amounts of Mo or the formation of a Mo-rich phase that is poorly crystallized. This would lead to lower peak intensity.

The XRD pattern for the Cu-Fe-Cr-Mo-B₄C alloy reveals the presence of Cu and (Fe, Ni) phases, similar to the Cu-Fe-Cr-Mo alloy, indicating that the primary crystal structure remains intact with the B₄C addition. As the XRD pattern does not detect any diffraction peaks for B₄C, it is likely due to its low quantity in the matrix or its dissolution or dispersion.

3.3. Microstructure of Fe Cores in Cu-Rich Regions

Figure 5 shows the Fe-rich region in Cu-Fe-Cr-Mo alloy and its surrounding microstructure. Figure 4 clearly shows the phase separation between the Cu-rich matrix and the Fe-rich region, as well as the Fe cores around the Fe-rich region.

Figure 5 illustrates that both Fe-rich regions and Fe cores are partitioned into dark grey and bright white areas. The bright white regions within the Fe-rich area display a continuous distribution, while in the Fe cores, the bright white regions are uniformly dispersed. The EDS results in

Table 3. Elemental distribution (at. %) in Cu-Fe-Cr-Mo alloy.

Points	Fe	Cu	Cr	Mo
A	70.00	6.70	17.67	5.64
B	7.10	89.94	1.93	1.03
C	68.89	4.10	18.88	8.13
D	64.43	0.92	21.77	12.88
E	74.51	7.94	16.23	1.31
F	6.27	91.37	2.02	0.34

show that the composition of points B and F in the Cu-rich matrix is almost all Cu, only a small amount of supersaturated Fe is dissolved in the Cu-rich matrix, and the content of Cr and Mo in the Cu-rich matrix is almost negligible. Based on the EDS outcomes at points A, C, and D, it appears that Fe, Cr, and Mo are the primary constituents of the white region. As the Mo level at A, C, and D is markedly higher than that of point E, it infers that Mo predominantly collects in the white area.

Given the presence of a very small amount of Mo in the Cu-rich matrix, most of the Mo is dispersed in the Fe-rich region, where brighter regions are formed in the Fe-rich region without phase separation from Fe and Cr. Despite the increase in the Mo content in these regions, the proportions of Fe and Cr, the two main components of RCF103, remain relatively close to those in Table 1.

SEM results of Cu-Fe-Cr-Mo-B₄C are shown in Figure 6a. In contrast to Cu-Fe-Cr-Mo, numerous small bright white aggregates were present in Cu-Fe-Cr-Mo-B₄C. Figure 6b depicts a selected bright white region, and its EDS surface scan outcomes are displayed in Figure 6c–h. The scans of the EDS surface reveal that there are clear separation characteristics between the Fe-rich and Cu-rich phases (as shown in Figure 6e,g). Additionally, the majority of Cr is located within the Fe-rich region (as shown in Figure 6h), which is comparable to the results from Figure 4 presenting the Cu-Fe-Cr-Mo and EDS findings. The prominent aspect to take notice of is the dispersion of Mo and B₄C. As per Figure 6f, the luminous white material in the Fe-abundant area exhibits a significant level of Mo, although the distribution profile of Mo only imprecisely coincides with the Fe-abundant area. Figure 6c,d illustrate that the distributions of C and B differ, suggesting that B₄C decomposes in the melt pool, and C assumes a distribution form distinct from B. Figure 6c evinces no noticeable zone of C aggregation, indicating that C should be evenly disseminated in the melt without the influence of other components. On the other hand, the region of aggregation for B coincides with the region of aggregation for Mo, suggesting the possibility of a Mo-B cluster or compound formation between B and Mo.

Unique Fe cores with a core–shell structure were also found in the Cu-Fe-Cr-Mo-B₄C alloy, as shown in Figure 7, and this structure of the Fe core has also been found in other studies [29].

4. Discussion

4.1. Mechanism of Formation of Macroscopic Cu-Fe Phase Separation

According to Figure 8, which represents the phase diagram of the Cu-Fe binary alloy, it can be deduced that the solidification process of the alloy follows a certain order. Initially, the dominant solidification phase is the native δ-Fe phase which first nucleates and then grows. Under the conditions of LC processing, the solidification interface generates a substantial amount of undercooling, leading to the local melting of the solid–liquid interface. As the alloy system solidifies, the overall free energy difference reduces with decreasing entropy, and the surface tension between the Cu-rich matrix and the Fe-rich region increases to minimize the surface energy [30]. The formation of Fe-rich regions, isolated and solidified within the Cu-rich matrix, is a direct result. The selected alloy for this study has a Cu atomic fraction of 54%~57%, as indicated by the phase diagram. Consequently, any liquid phase exhibiting supercooling greater than ΔTc (where the separation temperature for the liquid phase is lower than the miscibility gap) will decompose into the Fe-rich area and Cu-rich matrix. The findings in Figure 2 demonstrate that during the solidification of the alloy, liquid phase separation takes place, affirming that the extent of the actual undercooling necessary to induce such separation under the LC process conditions is beyond the corresponding undercooling threshold for liquid Fe cores. Due to the higher melting point of Fe than Cu, Fe has greater preferential nucleation kinetics, resulting in Fe cores nucleating predominantly and retaining a spherical shape.

It should be noted that the alloy compositions used here also contain certain components of Cr and Ni. For Cr, the same immiscible gap exists between Cr and Cu, and the EDS results herein show that Cr exists almost exclusively in the Fe-rich region, making the mechanism of liquid phase separation occurring between Cr and Cu somewhat similar to that between Fe and Cu. For Ni, Ni can change the temperature at which phase separation occurs in the alloy system, but the general pattern of phase separation is consistent. Therefore, the authors have made some simplifications to the composition in order to qualitatively discuss the liquid phase separation that occurs between Cu and Fe.

Furthermore, LC generates a melt pool with elevated transients and a rapid solidification rate [32,33]. This rate surpasses the necessary undercooling for liquid phase separation to occur in Cu-Fe-Cr alloys. In the melt pool, conspicuous macroscopic segregation is noticeable in all three alloys, reflecting differing flow rates between the Fe-rich region and the Cu-rich matrix. During solidification, the Cu-Fe-Cr alloy quickly enters the miscible gap and undergoes liquid phase separation.

When liquid phase separation takes place and the region rich in Fe starts to solidify, the allocation of solutes between the branches can be described in the following manner.

$$C_s^{*}=k_0{c}_0{\left(1-f_s\right)}^{\frac{k_0-1}{q}}$$

(1)

where

$$q=\left(1-\beta \right)\left(1-\frac{v}{\mu }\right),q>0$$

(2)

where k₀ is the equilibrium partition coefficient, c is the original concentration of the solute, f_s is the solid phase fraction, β is the solidification shrinkage, μ is the isothermal travel velocity, and v is the fluid flow velocity in the isothermal normal direction. Thus, the average solid-phase composition over a range can be derived as

$$\overline{C_s}={\displaystyle \int C_s^{*}d{f}_s}=\frac{k_0{c}_{0}q}{k_0-1+q}$$

(3)

When the initial composition c₀ of the Cu-Fe-Cr alloy solvate is constant, β is also constant. For q = 1, Equation (1) is the standard Scheil equation, indicating that no macroscopic segregation occurs in this case. Here,

$$\frac{v}{\mu }=-\frac{\beta }{1-\beta }$$

(4)

Previous studies have pointed out that the melt pool generated via LC has a very fast melt flow rate [34,35]. The flow rate of the liquid in the melt pool is tens of times higher than the laser scanning speed, far exceeding the critical flow rate for no macroscopic segregation to occur. Here, $\frac{v}{\mu }<-\frac{\beta }{1-\beta }$ and q > 1. According to the laser deposition simulation [34], the fluid in the melt pool is also observed to flow from above the melt pool along the centerline of the melt pool to the sides of the melt pool, while the solidification interface of the melt pool moves in the opposite direction. Consequently, the solute between the Cu-rich matrix and the Fe-rich region moves from the hot end to the cold end. Due to the higher melting point of Fe than that of Cu, the Fe-rich region excludes Cu as a solute towards both ends of the melt pool during solidification. This results in the macroscopic segregation illustrated in Figure 2.

4.2. Impact of Mo and B₄C on Macroscopic Phase Separation Characteristics

The inclusion of Mo and B₄C modifies the latent heat of crystallization of the Cu-Fe-Cr alloy, thus altering the mixing enthalpy and mixing entropy of the alloy system. Both mixing enthalpy and mixing entropy have intricate impacts on phase separation [36]. The outcomes of the thermodynamic calculation software JmatPro for the alloy system are presented in Figure 9. The formation of phases is determined by free energy minimization, rather than being independently determined by either the mixing entropy or the mixing enthalpy. Phase separation occurs in alloys with low mixing entropy and a positive mixing enthalpy [37], with the positive mixing enthalpy of the alloy composition being a significant contributor to the formation of the liquid miscibility gap. The enthalpies of atomic pairs’ mixing can be evaluated through the Miedema model, and they are connected to interatomic interactions. Whether the interatomic energy in the alloy is stronger or weaker than in the unmixed components at a particular molar fraction determines a tendency to form compounds (preferred neighborhoods of different atoms) or phase separations (preferred neighborhoods of the same atoms) [38]. According to Boer et al.’s [39] calculations, Fe-Cu atomic pairs have an enthalpy of mixing of +13 kJ/mol, which causes the Cu-Fe alloy to separate into a liquid phase. Furthermore, the mixing enthalpies of Fe-Cr and Cu-Cr atomic pairs are −1 kJ/mol and +12 kJ/mol, respectively, while the mixing enthalpies of Fe-Mo and Cu-Mo atomic pairs are −2 kJ/mol and +19 kJ/mol, respectively. This indicates that Cr and Mo have a greater tendency to be adjacent to Fe than to Cu, explaining why Mo can only be found in the Fe-rich region shown in Figure 5 and Figure 6.

If the melt pool is utilized as a reference system, it attains a quasi-steady state as the laser moves [40,41]. Although the temperature at every point within the melt pool remains constant during the quasi-steady state, heat conduction and convection still lead to heat and mass transfer inside the melt pool under the melt pool reference system. If a laser scan path point is employed as the reference system, it will undergo the creation, expansion, and solidification of the melt pool. At each point in the melt pool, the temperature and concentration change transiently with time, making it a non-equilibrium process. If the melt pool is viewed as a thermodynamic system, it will always tend to evolve towards a state of thermodynamic equilibrium. The addition of Mo and B₄C transforms the evolution of the melt pool into a state of thermodynamic equilibrium, as shown by the changes in enthalpy and entropy in Figure 9, and the change in this process is objectively reflected.

The calculations displayed in Figure 9a demonstrate a reduction in the mixing enthalpy of the alloy during the solidification phase. The Cu-Fe-Cr-Mo alloy exhibits its lowest mixing enthalpy during the initial solidification, implying a detrimental influence of Mo on the phase separation in this stage. This may be attributable to the distribution of Mo already nucleated within the melt pool, impeding the transport of other components. The arrow marked with the forbidden symbol in Figure 10a illustrates this effect. As solidification continues, the mixing enthalpy of the Cu-Fe-Cr alloy decreases in comparison to that of the Cu-Fe-Cr-Mo alloy. This suggests that the Mo, which nucleates in the late stage of solidification, begins to promote phase separation through a stirring effect. This stirring effect is reflected in the fluid velocity within the melt pool, which approximates Darcy’s law,

$$v=\frac{K\left(\nabla p+\rho g\right)}{\eta f_s},$$

(5)

where K is the permeability coefficient, which is related to the dendrite structure and the dendrite gap. P is the pressure acting on the interdendritic fluid, ρ is the fluid density, g is the gravitational acceleration, and η is the fluid viscosity. Obviously, the flow velocity of the fluid in the melt pool is related to the density of the melt. When molybdenum is introduced to the Cu-Fe-Cr alloy, it solidifies quickly and moves rapidly through the molten pool due to its high melting point and density. This stirring effect has a significant impact on the Fe-rich region and is evident in Figure 2b, which shows the presence of continuous Fe-rich regions and numerous free Fe cores within the Cu-rich matrix.

The alteration in mixing enthalpy of the Cu-Fe-Cr alloy and the Cu-Fe-Cr-Mo alloy implies that the impact of Mo on phase separation is intricate. On the one hand, the pre-existing nucleated Mo impedes the merger of homogeneous phases during phase separation, leading to a widespread scatter of Fe in the matrix that is rich in Cu and make them challenging to amalgamate. However, once the stirring effect of Mo driven by Marangoni convection in the melt pool becomes apparent, the pool has already progressed to the latter stages of solidification. Consequently, numerous Fe cores are unable to undergo adequate collision and fusion, resulting in a considerable quantity of free Fe cores that ultimately become enclosed by the Cu-rich matrix. The enthalpy mixing of the Cu-Fe-Cr-Mo-B₄C alloy is notably greater than that of the other two compositions during nearly the entire solidification process. Furthermore, the Cu-rich matrix in Figure 2c has fewer Fe cores and more consistent Fe-rich regions, which strongly suggests that the inclusion of B₄C is a significant factor in the phase separation phenomenon.

The entropy of the Cu-Fe-Cr-Mo alloy during solidification is less than that of the Cu-Fe-Cr alloy, suggesting that the inclusion of Mo diminishes the alloy system’s latent heat of crystallization. Since entropy assesses the degree of disorder, a decline in the mixing entropy signifies a reduction in the alloy system’s degree of disorder. This indicates that the introduction of Mo has an unfavorable impact on the separation of the liquid phase, demonstrated by the presence of numerous independent Fe cores surrounding the Fe-rich area, as depicted in Figure 2b. When a large number of Fe cores are yet to collide and merge with each other, the melt temperature is likely to decline to the melting point of Cu under severe undercooling. In contrast to this, Figure 2a demonstrates a significant decrease in the number of Fe cores in the Cu-rich matrix.

The inclusion of B₄C results in a notable rise in the entropy of the alloy system. The data presented in Figure 9b illustrate that the entropy of Cu-Fe-Cr-Mo-B₄C is dissimilar to that of the Cu-Fe-Cr-Mo and Cu-Fe-Cr alloys until the start of the nucleation process in the Cu-rich matrix. The augmented mixing entropy displays an increased level of disorder and superior miscibility, as showcased by the more uniformly dispersed Fe-rich regions on the Cu-rich matrix depicted in Figure 2c. Based on previous research, it has been found that elements C and B widen the miscibility gap between Fe and Cu, as shown in Figure 10b [42,43]. This extended duration of the melt pool within the immiscible gap facilitates liquid phase separation [44,45]. Although the preferentially shaped Mo nuclei impede the collision and fusion of Fe cores during the initial phase of phase separation, the addition of B and C expands the miscibility gap of the melt, resulting in a gradual weakening of the hindering effect of Mo as phase separation progresses. Due to the constant mutual attraction between the Fe cores, the number of Fe particles obstructed by Mo increases. At a certain point, the mutual attraction of Fe particles exceeds the obstruction of Mo. At this time, Fe breaks through the Mo obstruction and collides with other Fe cores. Moreover, the miscibility gap expanded by elements C and B provides ample time for the Fe cores to impact and combine with each other, ultimately leading to a noteworthy reduction in the quantity of unbound Fe cores within the Cu-rich matrix. This alteration also causes a shift in the distribution of Fe core particles to smaller sizes (Figure 3e), which facilitates the collision and fusion of the Fe cores. Nevertheless, it is evident that the appearance of adjoining Fe-rich regions in Figure 2b contrasts significantly with that of Figure 2c. In the Cu-Fe-Cr alloy, phase separation almost follows the process illustrated in Figure 8. The separation primarily produces continuous Fe-rich areas dispersed throughout the Cu-rich matrix, as seen in Figure 11a. Mo has been noted to break up the Fe nuclei in the pre-phase separation stage. However, even though the miscibility gap remains the same as in the Cu-Fe-Cr alloy, the final solidified Fe-rich region displays a shape that breaks up the Fe-rich area based on the Fe-rich region in Figure 2a, as depicted in Figure 11b.

When B₄C widens the miscibility gap, the Fe cores can collide and fuse. However, the influence of Mo has driven the Fe cores to disperse uniformly in the melt. As a result, the Fe cores are more likely to fuse with their closest other Fe core (see Figure 11c), leading to the formation of Fe-rich regions with a diffuse-like appearance. Furthermore, the Fe-rich area is extensively elongated into a diffuse shape as demonstrated in Figure 11d. B₄C particles act as heterogeneous nucleation sites, contributing to a more homogeneous distribution of Fe-rich droplets in the Cu-rich matrix. During rapid solidification, these nucleation sites hinder the uncontrolled growth and coalescence of Fe cores, thus allowing Fe-rich droplets to maintain a more rounded shape due to the minimization of surface free energy. In addition, B₄C has a high affinity for Fe and can modify the local concentration of Fe as well as the surface tension effect in the molten alloy, resulting in the rounded morphology of the Fe-rich region in Figure 2c.

It is worth noting that Fe cores can still be found freely in the Fe-rich region, as depicted in Figure 2c. This is mainly attributed to the rapid solidification of the melt pool induced by laser cladding. However, the addition of B₄C recapacitates the Fe cores for collision and fusion, but this assumes that it is carried out in the liquid phase. Once the solid–liquid interface of the melt pool has passed through the Fe-rich region, only the unfused free Fe cores can exist in the Cu-rich matrix. However, it is evident that a significant portion of the Fe phase remains fused into substantial Fe-rich regions, notwithstanding the rapid solidification rate of the melt pool.

4.3. Alteration of Fe Cores Microstructure by Mo and B₄C

Given that Mo has a higher melting point than Fe, Cr, and Cu, and its solubility in Cu-rich melt is significantly lower than in Fe-rich melt, Mo first nucleates in Fe-rich melt during solidification. The results in Figure 2b already indicate that Mo increases the number of Fe cores, hence separated Fe cores will unavoidably take away some of the previously formed Mo once Fe cores separate from the Fe-rich region. When the Fe core enters the region rich in Cu, its cooling rate hastens because of the higher thermal conductivity of Cu compared to that of Fe. Consequently, there is a rapid transfer of heat from the free Fe core to the Cu-rich matrix, resulting in a much faster solidification rate of the Fe core than that of the Fe-rich area, creating the white area in the Fe core as illustrated in Figure 5b. In the Fe-rich region, as a consequence of the slower solidification rate, the nucleated Mo has adequate time to collide and fuse with each other, thereby resulting in the continuous white Mo-rich region as depicted in Figure 5a. In contrast to point E, Mo is predominantly present within the white dendrites and forms a eutectic organization with Fe-Cr as depicted in Figure 5b. The EDS findings at point E in Figure 5a suggest that a specific extent of supersaturated Cu is present in the Fe-rich area. This region is alike in its element distribution to the Fe core surrounding it (point A), demonstrating the presence of supersaturated Cu in the core. The possibility is raised that the core has undergone a secondary separation from its internal supersaturated copper. The emergence of secondary liquid–liquid phase separation primarily results from the strong undercooling of LC. Such undercooling causes the initial formation of liquid Fe cores which quickly enter the miscibility gap and thus trigger a simultaneous, spontaneous second liquid–liquid phase separation within the primary phase-separated Fe cores [35]. A certain degree of Fe in a supersaturated state (B and F points) is also observable in the Cu-rich matrix, whereas the Cu-rich matrix, which holds almost no Cr or Mo, suggests that the solid solution of both in Cu is close to zero.

In order to investigate whether the bright white part in Figure 6 is just a solid solution of Mo and B or an in situ synthesized compound, the possible reactions of Mo and B₄C in the melt pool were investigated. B₄C absorbs a large amount of laser beam radiation and decomposes into elements C and B. According to the theory of chemical thermodynamics, spontaneous chemical reactions occur when the Gibbs free energy has a negative value [45]. The Gibbs free energy changes at different temperatures as

$$\mathsf{\Delta }G=\mathsf{\Delta }H-T\mathsf{\Delta }S$$

(6)

where ΔG is the Gibbs free energy of chemical reaction; ΔH is the enthalpy change of chemical reaction; T is the temperature; and ΔS is the entropy change of the chemical reaction. The following reactions may occur between Mo and B₄C:

$$\mathrm{Mo}+{\mathrm{B}}_4\mathrm{C}\to \mathrm{M}{\mathrm{o}}_2\mathrm{B}+\mathrm{M}{\mathrm{o}}_2\mathrm{C}$$

(7)

$${\mathrm{B}}_4\mathrm{C}\to \mathrm{B}+\mathrm{C}$$

(8)

$$\mathrm{M}\mathrm{o}+\mathrm{B}\to \mathrm{M}{\mathrm{o}}_2\mathrm{B}$$

(9)

$$\mathrm{Mo}+\mathrm{C}\to \mathrm{M}{\mathrm{o}}_2\mathrm{C}$$

(10)

According to the calculation, the Gibbs free energy change of the above reactions is shown in Figure 12a. The temperature at each point inside the molten pool decreases with time during solidification, and the equilibrium constant K is a measure of the degree of chemical reaction with time. The greater the K value of the chemical reaction, the greater the concentration of the product at equilibrium [46]. In order to analyze the most likely compounds produced in laser melting, it is necessary to calculate K. Based on the Gibbs free energy calculation, the reaction equilibrium constant K can be calculated as

$$K=\frac{{[A_1]}^{s_1}\cdot {[A_2]}^{s_2}\cdot \dots \cdot {[A_n]}^{s_n}}{{[B_1]}^{s_1}\cdot {[B_2]}^{s_2}\cdot \dots \cdot {[B_m]}^{s_m}}$$

(11)

where ${[A_n]}^{s_n}$ is the activity or partial pressure and chemical coefficient of the product, respectively, and ${[B_m]}^{s_m}$ is the activity or partial pressure and chemical coefficient of the reactant, respectively. Thus,

$$\ln K=\frac{-\mathsf{\Delta }G}{RT}$$

(12)

where R is the gas constant and T is the reaction temperature. By computing the equilibrium constants of the aforementioned likely reactions, the findings exhibited in Figure 12b can be attained. The lnK ratings for Reaction (4) are positive over the entire melt pool temperature range, indicating that Reaction (4) can proceed spontaneously, provided that Reaction (2) occurs. Reaction (2) takes place at melt pool temperatures above 3000 K. Nevertheless, the absence of Mo₂C in the Cu-Fe-Cr-Mo-B₄C alloy has been confirmed by the EDS results, indicating that Reaction (2) and therefore Reaction (4) do not happen. Moreover, the lnK values for both Reaction (1) and Reaction (3) are negative, suggesting that Reactions (1) and (3) do not occur within the temperature range of the melt pool. Furthermore, according to the predicted enthalpies of formation of binary molybdenum compounds by Boer et al. [37], the interatomic energies between the various forms of Mo-B binary compounds are negative, indicating that there is indeed a tendency for Mo and B to be attracted to each other and form Mo-B clusters, combining with the XRD results of Figure 4.

During the initial phase of the liquid phase separation process, the powerful convection that occurs within the molten pool results in the prompt development of separate Fe cores enclosed in a continuous matrix that is rich in Cu. As Fe cores grow larger, the surrounding concentration field causes significant drifting and collisions between the neighboring Fe cores. Additionally, as the Marangoni motion acts against the temperature gradient, it causes further collisions between the remaining Fe cores and becomes increasingly important. Both B and C will enhance the undercooling of the melt, causing a constant change in the concentration of the surrounding Cu-rich matrix, and the attractive interactions between the Fe cores will further enhance droplet collisions and collision induction, leading to the formation of larger, faster-moving Fe cores driven by the Marangoni effect. When the increase in undercooling makes the thermal Marangoni migration faster than the solute Marangoni migration, the Cu-rich matrix exhibits an inward radial migration, creating an encircling potential for the Fe core and eventually forming the unique core–shell structure shown in Figure 7.

4.4. Microhardness

Figure 13 shows the microhardness results for the Cu-Fe-Cr composite. The measurement range starts from the Fe-rich region and ends at the Cu-rich matrix. Similar to Dai’s study [47], a sharp decrease can be found in Figure 13, which suggests the existence of a boundary between the Fe-rich region and the Cu-rich matrix. The results show that the Cu-Fe-Cr alloy has the lowest microhardness in both the Fe-rich zone and the Cu-rich matrix. After the addition of Mo, the microhardness of Cu-Fe-Cr-Mo alloys increased by approximately 100 HV in the Fe-rich area when compared to Cu-Fe-Cr alloys. Furthermore, the microhardness of the Cu-rich matrix is nearly equivalent to twice that of Cu-Fe-Cr alloys. The strengthening mechanism of molybdenum varies between areas enriched in iron and substrates enriched in copper. The findings in Table 3 indicate that molybdenum is primarily located within the Cr-Mo eutectic, distributed evenly in the iron-rich domains to create a dispersed solid solution reinforcement. Molybdenum is not present in the copper-rich matrix but acts as a tissue refiner and enhances the matrix strength. The addition of B₄C further increased the microhardness of the Fe-rich region, resulting in the Fe-rich region exceeding 600 HV, indicating that the strengthening effect of B₄C on the Fe-rich region of F is obvious [48]. When B₄C is added, it acts as a hard phase reinforcement. Dispersed B₄C particles have the potential to effectively enhance Fe-rich areas, akin to precipitation hardening, while hard B₄C particles can hinder the dislocation motion, thus increasing the difficulty for dislocations to move through the material. Nonetheless, the strength of the Cu-rich matrix stays almost identical to that of the Cu-Fe-Cr-Mo alloy, signaling that the B₄C particles are influenced by the phase separation and the strengthening of the Cu-rich matrix is challenging.

Table 4. Microhardness comparison (the microhardness values are averages of multiple measurements).

Cu-Fe Systems	Microhardness(HV) in Fe-Rich Region	Microhardness(HV) in Cu-Rich Region
Cu65Fe35 [20]	438	153
Cu61.1Fe32.9Si6 [16]	594	243
Cu-Fe-Cr	425	126
Cu-Fe-Cr-Mo	553	187
Cu-Fe-Cr-Mo-B4C	668	183

presents additional microhardness measurements from related studies. Notably, the mechanical properties of Cu-Fe alloy systems are significantly impacted by the relative content of Cu and Fe. However, due to the varying relative contents of Cu and Fe in current studies, making accurate comparisons is challenging. As a result, the properties of different Cu-Fe alloy systems can only be measured qualitatively.

Table 2 results indicate that without reinforcing phase incorporation, the microhardness of Cu-Fe-Cr alloys obtained in this study is comparable to that reported in other studies. However, Ref. [16] demonstrates that the Cu-rich matrix has notably higher hardness, implying a potential strengthening effect of Si on the Cu-rich matrix.

4.5. Wear Performance and Wear Mechanism

Figure 14 shows the results of the coefficient of friction for Cu-Fe-Cr alloys of different compositions. The calculated average coefficients of friction for 20–30 min of wear at the smooth wear stage are 0.67 for the Cu-Fe-Cr alloy, 0.6 for the Cu-Fe-Cr-Mo alloy, and 0.54 for the Cu-Fe-Cr-Mo-B₄C alloy. During the early stages of wear, the friction coefficient of the Cu-Fe-Cr alloy exceeds that of the Cu-Fe-Cr-Mo alloy and Cu-Fe-Cr-Mo-B₄C alloy. However, as the wear continues, the coefficient of friction of the Cu-Fe-Cr-Mo-B₄C alloy shifts from leading to lagging behind the Cu-Fe-Cr-Mo alloy, signifying a change in the type of friction. It is noteworthy that the curve for the Cu-Fe-Cr-Mo alloy depicts a more pronounced dip at the 13th second compared to the other components, suggesting an abrupt alteration in the friction coefficient at that moment. This phenomenon is less discernible in the curves of the remaining two alloys. Combined with Figure 2, it is suggested that the irregular distribution of the Cu-rich matrix and Fe-rich region may be the cause. During this time, the pair of grinding balls may have coincidentally moved to the Cu-rich matrix with a lower hardness. Additionally, there are fewer discrete Fe nuclei distributed within the Cu-rich matrix in this region compared to others. If more uniformly distributed Cu-Fe-based alloys are to be achieved, these findings indicate that additional measures favoring miscibility, such as electromagnetic stirring, need to be implemented. Preliminary test results showed that the chemical composition of the Cu-Fe-Cr alloy had an effect on the friction properties, which was also confirmed in other studies [29].

The comparative results in

Table 5. Microhardness comparison (the values of the friction coefficients of Cu₈₈Fe₁₂ are estimates based on the results of friction experiments).

Cu-Fe Systems	Friction Coefficient
Cu88Fe12 [29]	~0.8
Cu86.5Ni7.4Si3.8Fe1.1 [49]	0.73
Cu-Fe-Cr	0.67
Cu-Fe-Cr-Mo	0.6
Cu-Fe-Cr-Mo-B4C	0.54

show that the addition of Cr, Mo, and B₄C has significantly enhanced the wear resistance of immiscible Cu-Fe alloys.

Figure 15 depicts the surface wear morphology of the Cu-Fe-Cr composite alloy. The predominant forms of wear encompass plastic deformation and plowing. Notably, the Cu-Fe-Cr alloy exhibits the deepest groove depth, amounting to roughly 43 μm, while the Cu-Fe-Cr-Mo alloy manifests a diminished groove depth of 18 μm. Additionally, the Cu-Fe-Cr-Mo-B₄C alloy features a further reduced groove width of 14 μm, indicating its superiority in wear resistance amongst these alloy variants. The EDS analysis of the trench base indicates the presence of a substantial amount of iron. As depicted in Figure 2, the macroscopic surface morphology reveals that the Fe-rich area is elevated compared to the Cu-rich matrix, thereby suggesting that the loss incurred by the Cu-rich matrix is greater than that of the Fe-rich region during the corrosive process. When the opposing grinding ball contacts the surface, it strips away the Fe-rich region, resulting in the ground iron chips being integrated into the Cu-rich matrix. The addition of Mo enhances the wear resistance of the alloys. Figure 15b,e show significantly improved groove smoothing compared to the Cu-Fe-Cr alloy. However, regrettably, wear-induced cracks were observed in the grooves, which suggests a relationship with compositional inhomogeneities. Figure 9b shows that the inclusion of Mo reduces the mixing entropy of the alloy and impedes the collision and fusion of Fe nuclei that have been preferentially nucleated, leading to a significant amount of Fe cores remaining within the Cu-rich matrix. It has been previously established that the pair of grinding balls will have a higher likelihood to come into contact with the Fe-rich area. Consequently, if a significant amount of Fe cores is amassed in a particular region within the Cu-rich matrix, then there is a possibility that the pair of grinding balls will strip the Fe cores simultaneously. In such a scenario, as Cu is considerably less hard than Fe, cracks will arise when the Fe is detached from the Cu. For Cu-Fe-Cr alloys, the Fe-rich region is widely dispersed throughout the Cu-rich matrix, despite the significant difference in hardness between the two. Therefore, if there is any loss of the Cu or Fe phase, it can only reattach to the corresponding phase surface. This process makes it difficult to create scratches on the surface after it is subjected to rubbing. The chance of cracking is reduced by the Cu-Fe-Cr-Mo-B₄C alloy due to the increased mixing entropy caused by the addition of B₄C. This results in a more uniform distribution of the Cu-rich matrix and Fe-rich regions after liquid phase separation.

The wear resistance of the Cu-Fe-Cr-Mo-B₄C alloy was found to be superior in comparison to that of the Cu-Fe-Cr alloy, as the results depicted in Figure 16 suggest. This aligns with the findings presented in Figure 14. Figure 17 expounds the underlying principle behind this observation. Notwithstanding the large Fe-rich regions characterizing the Cu-Fe-Cr and Cu-Fe-Cr-Mo alloys, distinctly different wear forms appear. Specifically, in the Cu-Fe-Cr alloy, the hardness of the Fe-rich region is 2.5 times more than that of the Cu-rich matrix. The Fe-rich area comprises approximately 15% Cr, thereby increasing its brittleness. The insertion of grinding ball pairs within this zone generates iron flakes, as evidenced in Figure 17a. These flakes possess distinctive edges and, as the grinding ball moves, the iron chips act as abrasive particles, which consequently deepens the grooves. This, in turn, results in the grooves of Cu-Fe-Cr alloy reaching a significant depth, depicted in Figure 15a. The incorporation of Mo enhances the hardness of the Cu-rich matrix and to some extent mitigates the damage to the Cu-rich matrix caused by the flake iron. Conversely, the abundance of independent Fe cores in the Cu-rich matrix blunts the originally relatively sharp Fe chips (Figure 17b), which also reduces the destructive power of the Fe chips on the Cu-rich matrix.

The incorporation of B₄C significantly enhances the wear resistance of the alloy caused by a blend of factors. Figure 13 results demonstrate that, as opposed to the Cu-Fe-Cr-Mo alloy, where there is a greater hardness contrast between the Cu-rich matrix and the Fe-rich region, the introduction of B₄C minimizes the hardness variation between the two regions. This leads to the absence of cracking comparable to Figure 15b and results in flatter grooves after friction. Furthermore, the addition of B₄C decreases the quantity of Fe cores not embedded in the Cu-rich matrix. Even if the Fe rich matrix is initially damaged by the grinding ball, the stripped iron fragments are ground into the Cu-rich matrix, ultimately producing non-sharp chips similar to those depicted in Figure 17b. As a result, the destruction of the softer Cu-rich matrix is slowed down.

4.6. Magnetization Properties

Figure 18 illustrates the hysteresis loop detection outcomes of the Cu-Fe-Cr composite alloy at room temperature (296.15K). Among these alloys, Cu-Fe-Cr-Mo-B₄C exhibits the strongest saturation magnetization (124 emu/g) and coercivity (90 Oe), while Cu-Fe-Cr displays the weakest remanence (2 emu/g) and initial permeability. The saturation magnetization of Cu-Fe-Cr-Mo-B₄C is higher compared to the Cu-Fe-Ni-Cr-Si alloy, which was prepared by Dai (96.2 emu/g) [47], and the Cu-Fe-Cr-Si-C alloy, which was prepared by Zhou (84.6 emu/g) [17], as shown in

Table 6. Magnetization properties comparison.

Cu-Fe Systems	Saturation Magnetization (emu/g)	Coercivity (Oe)
Cu-Fe-Cr-Si-C [17]	84.6	0.65
Cu-Fe-Ni-Cr-Si [47]	96.2	7.24
Cu-Fe-Cr	68	68
Cu-Fe-Cr-Mo	118	55
Cu-Fe-Cr-Mo-B4C	124	90

.

The initial permeability is influenced by factors such as the magnetic grain size, orientation, and denseness. Sequentially, from the Cu-Fe-Cr alloy to the Cu-Fe-Cr-Mo-B₄C alloy, the Fe core size decreases, while the saturation magnetization intensity increases, resulting in an increase in the initial permeability. The initial permeability is influenced by factors such as the magnetic grain size, orientation, and denseness. The grain size variation significantly affects the coercivity. For approximately spherical grains, the coercivity increases as the grain size decreases, reaching a peak value, before decreasing further with the diminishing grain size. As the Cu-Fe-Cr-Mo-B₄C alloy has the smallest average grain size yet large coercivity, it implies that its average grain size is smaller than the size of a single domain at this stage. Simultaneously, the Cu-Fe-Cr-Mo-B₄C alloy amplifies the irreversible movement and rotation of the domain wall owing to the heightened internal stress, thereby leading to increased coercivity. The hysteresis lines depicted in Figure 18 signify that the three alloys are not transiently magnetic substances, implying the absence of internal antimagnetization nuclei in the alloys. Compared to Cu-Fe-Cr alloys, Cu-Fe-Cr-Mo and Cu-Fe-Cr-Mo-B₄C alloys possess relatively smaller Fe cores. As a result, the intensity of the demagnetization field generated around the Fe cores during the magnetization process is decreased, thereby leading to an elevated level of remanent magnetization. Compared to the Cu-Fe-Cr-Mo alloy, the Cu-Fe-Cr-Mo-B₄C alloy contains a greater number of small Fe cores. As a result, the intensity of the demagnetization field around the Fe cores is decreased, leading to a higher remanence than that of the Cu-Fe-Cr-Mo alloy. However, the remanent magnetization of Cu-Fe-Cr-Mo-B₄C is actually lower than that of Cu-Fe-Cr-Mo. This is most likely because the addition of B₄C leads to a large number of Fe cores fusing into an Fe-rich region (Figure 2c). This consequently increases the remanent magnetic field strength, which should ideally be lower, resulting in a lower remanent magnetization.

The findings presented in Table 6 demonstrate a notable discrepancy between the alloy system obtained in this investigation and previous research in terms of coercivity. Coercivity reflects the magnetic field strength necessary to fully demagnetize an alloy once it has been magnetized. The coercivity of the alloy system developed in this study surpasses that of prior studies. This points to the weakness of the alloy concerning its demagnetization characteristics. This is an area that will be focused on to enhance in the future.

5. Conclusions

The present paper elucidates the microstructure and coating properties of Cu-Fe-Cr composite coatings with varied chemical compositions. The results of this study are as follows:

• The composition of the alloy influences liquid phase separation. Mo obstructs the collision and fusion of Fe cores, resulting in an uneven distribution of composition. The miscibility is enhanced by B₄C, reducing property discrepancies due to macroscopic segregation.
• Molybdenum in Cu-Fe-Cr-Mo alloys primarily concentrates on the branched crystals created by chromium, forming a eutectic structure with it. Despite there being no chemical reaction between B₄C and Mo in the Cu-Fe-Cr-Mo-B₄C alloy, due to the decomposition of B₄C and the negative interatomic energy between B and Mo, Mo and B tend to attract each other and are likely to form Mo-B clusters.
• The alloy of Cu-Fe-Cr-Mo exhibits the greatest micro-hardness, measuring approximately 600 HV, whereas Cu-Fe-Cr-Mo-B₄C displays the superior wear resistance, with a friction coefficient of 0.54.
• Compared to the Cu-Fe-Cr alloy, the Cu-Fe-Cr-Mo alloy and Cu-Fe-Cr-Mo-B₄C exhibit stronger saturation magnetization (118 emu/g and 124 emu/g) and lower coercivity (90Oe and 55Oe), indicating their potential for application in the field of magnetic materials.

While LC shows promising prospects as a speedy solidification process for creating Cu-Fe alloys, uncertainties in the LC method still persist. The present paper studies Cu-Fe alloys with different chemical compositions only on the basis of fixed process parameters, which also have an important influence on the properties of Cu-Fe alloys, which is the focus of the authors’ future research.